High carbon hot-rolled steel sheet and method for manufacturing the same

ABSTRACT

The high carbon hot-rolled steel sheet contains, in terms of percentages of mass, 0.10 to 0.7% C, 2.0% or less Si, 0.20 to 2.0% Mn, 0.03% or less P, 0.03% or less S, 0.1% or less Sol.Al, 0.01% or less N, and the balance being Fe and inevitable impurities, and has a structure of ferrite having 6 μm or less average grain size and carbide having 0.10 μm or more and less than 1.2 μm of average grain size. The volume ratio of the carbide having 2.0 μm or more of grain size is 10% or less. The volume ratio of the ferrite containing no carbide is 5% or less. The manufacturing method thereof has the steps of hot-rolling, primary cooling, holding, coiling, acid washing, and annealing. The primary cooling step is to cool the hot-rolled steel sheet down to cooling termination temperatures ranging from 450° C. to 600° C. at cooling rates of higher than  120 ° C./sec. The holding step is to apply secondary cooling to hold the primarily cooled hot-rolled steel sheet at a temperature range from 450° C. to 650° C. until coiling.

FIELD OF THE INVENTION

The present invention relates to a high carbon hot-rolled steel sheethaving excellent ductility and stretch-flange formability, and amanufacturing method thereof.

DESCRIPTION OF THE RELATED ARTS

High carbon steel sheets employed for tools, automobile parts (gear,transmission), and the like are subjected to heat treatment such asquenching and tempering after punching and forming thereof. The requestsof users who conduct the working on these components include improvementin bore expansion (burring) property in the forming process afterpunching, as well as the elongation characteristic which is an index ofductility for forming the steel sheet into complex shapes. The burringproperty is evaluated by the stretch-flange formability as one ofpress-forming properties. Consequently, there are wanted the materialshaving excellent stretch-flange formability as well as ductility.

Regarding the improvement in the stretch-flange formability of highcarbon steel sheets, several technologies have been studied. Forexample, JP-A-11-269552 and JP-A-11-269553, (the term “JP-A” referred toherein signifies the “Japanese Patent Laid-Open Publication”), disclosea method for manufacturing medium to high carbon steel sheets havingexcellent stretch-flange formability in a process after cold-rolling.The disclosed technology employs a hot-rolled steel which contains 0.1to 0.8% C by mass, having a metallic structure substantially consistingof ferrite phase and pearlite phase, having, at need, the area rate ofproeutectoid ferrite of at or higher value determined by the C content(% by mass), and having 0.1 μm or larger distance between pearlitelamellas. To the hot-rolled steel sheet, cold-rolling is given by 15% orhigher rolling rate, followed by three-stage or two-stage annealingwhile holding the steel sheet in three steps or two steps of temperatureranges for a long time.

Also JP-A-2003-13145 discloses a method for manufacturing a high carbonsteel sheet having excellent stretch-flange formability, which contains0.2 to 0.7% C by mass, has average grain size of carbide in a range from0.1 to 1.2 μm, and has a volume ratio of carbide-free ferrite grains of10% or less. The disclosed technology is a process in which thehot-rolling is given at finishing temperatures of (Ar₃ transformationpoint −20° C.) or above, the cooling is given to cooling terminationtemperatures of 650° C. or below at cooling rates of higher than 120°C./sec, the coiling is given at temperatures of 600° C. or below, theacid washing is given, and then the annealing is given at annealingtemperatures ranging from 640° C. to Ac₁ transformation point.

According to the technologies disclosed in JP-A-11-269552 andJP-A-11-269553, the ferrite structure is made by the proeutectoidferrite and does not include carbide. As a result, the stretch-flangeformability is not necessarily favorable, though the material is softand shows excellent ductility. A presumable reason of the phenomenon isthe following. During punching the steel sheet, the area of proeutectoidferrite significantly deforms in the vicinity of a punched end face,which induces significant difference between the deformation of theproeutectoid ferrite and that of the ferrite containing spheroidcarbide. As a result, stress concentrates to the peripheral zones ofgrain boundary where the deformation significantly differs therebetween,thereby generating voids at interface between the spheroid structure andthe ferrite. Since the voids grow to cracks, the stretch-flangeformability is ultimately deteriorated.

A countermeasure to the phenomenon may be the one to apply strengthenedspheroidizing annealing, thereby softening the entire structure. In thismeasure, however, the spheroidized carbide becomes coarse to become theorigin of void during the forming step, and the carbide becomes lesssoluble in the heat treatment step after the forming to cause thedecrease in quenched strength.

Furthermore, recent requirement for the forming level has increased morethan ever from the point of increase in productivity. Consequently,burring in high carbon steel sheet also likely induces crack generationat punched end face caused by the advanced level of forming. Therefore,the high carbon steel sheets are also requested to have highstretch-flange formability.

In this regard, the inventors of the present invention developed atechnology disclosed in JP-A-2003-13145 aiming to provide a high carbonsteel sheet having excellent stretch-flange formability and inducingvery few cracks at punched end face, which steel sheet is manufacturedwithout applying time-consuming multi-stage annealing. The technologyallowed manufacturing a high carbon hot-rolled steel sheet havingexcellent stretch-flange formability.

On the other hand, recent uses of driving system components and the likerequest increased strength also in the non-heat treating parts,specifically in integrally formed components for attaining higherdurability and lighter weight, thus the steel sheets as the startingmaterial are requested to have 440 MPa or higher tensile strength (TS).That kind of request with the aim to reduce the manufacturing cost ofcomponents has led a request to supply hot-rolled steel sheets.

The integral forming process has more than ten pressing steps, and isconducted in a complex combination of forming modes including not onlyburring but also stretching and bending. Accordingly, the integralforming has faced the simultaneous requests of stretch-flangeability andelongation.

According to the technology disclosed in JP-A-2003-13145, however,achieving TS=440 MPa (73 point or more as HRB hardness) not necessarilyattains satisfactory stretch-flange formability. That is, the technologycannot satisfy stably the requirements of both that level of TS and thestretch-flange formability. Furthermore, the disclosed technology doesnot refer to the elongation.

Adding to the above technology, the technology disclosed inJP-A-2003-13145 generates transformation heat after cooling, whichincreases the temperature to enhance the precipitation of proeutectoidferrite and the pearlite transformation, thereby inducing growth ofcoarse carbide and uneven carbide distribution to likely deteriorate thecharacteristics.

SUMMARY OF THE INVENTION

It is an object of the present invention to provide a high carbonhot-rolled steel sheet having 440 MPa or higher tensile strength andgiving excellent ductility and stretch-flange formability, generatingvery few cracks at punched end face, and which steel sheet can bemanufactured without applying time-consuming multi-stage annealing.

The inventors of the present invention conduced intensive studies on theeffect of components and microscopic structures of high carbon steelsheet on ductility and stretch-flange formability while securingstrength thereof, and found that the ductility and the stretch-flangeformability of steel sheet are significantly affected by not only thecomposition of the steel, the shape and quantity of carbide, but alsothe dispersed state of carbide. That is, it was found that the ductilityand the stretch-flange formability of high carbon hot-rolled steel sheetare improved by controlling each of the carbide shape in terms ofaverage grain size of carbide and volume ratio of carbide having 2.0 μmor larger grain size, and the dispersed state of carbide in terms ofvolume ratio of carbide-free ferrite grains and average grain size offerrite.

The present invention provides a high carbon hot-rolled steel sheetconsisting essentially of, in terms of percentages of mass, 0.10 to 0.7%C, 2.0% or less Si, 0.20 to 2.0% Mn, 0.03% or less P, 0.03% or less S,0.1% or less Sol.Al, 0.01% or less N, and balance of Fe and inevitableimpurities, and having a structure of ferrite having 6 μm or smalleraverage grain size and carbide having 0.10 μm or larger and smaller than1.2 g m of average grain size. The volume ratio of the carbide having2.0 μm or larger grain size is 10% or less, and the volume ratio of theferrite containing no carbide is 5% or less. The high carbon steel sheetgives excellent ductility and stretch-flange formability.

The high carbon hot-rolled steel sheet may further contain at least oneelement selected from the group consisting of, in terms of percentagesof mass, 0.05 to 1.5% Cr and 0.01 to 0.5% Mo.

The high carbon hot-rolled steel sheet may further contain at least oneelement selected from the group consisting of, in terms of percentagesof mass, 0.005% or less B, 1.0% or less Cu, 1.0% or less Ni, and 0.5% orless W.

The high carbon hot-rolled steel sheet may further contain at least oneelement selected from the group consisting of, in terms of percentagesof mass, 0.05 to 1.5% Cr and 0.01 to 0.5% Mo, and further at least oneelement selected from the group consisting of, in terms of percentagesof mass, 0.005% or less B, 1.0% or less Cu, 1.0% or less Ni, and 0.5% orless W.

The high carbon hot-rolled steel sheet may further contain at least oneelement selected from the group consisting of, in terms of percentagesof mass, 0.5% or less Ti, 0.5% or less Nb, 0.5% or less V, and 0.5% orless Zr.

The content of Si is preferably from 0.005 to 2.0% by mass. From thepoint of securing strength after annealing, the Si content is morepreferably 0.02% or more. From the point of surface property, the Sicontent is more preferably 0.5% or less.

The content of Mn is preferably from 0.2 to 1.0% by mass.

A preferable range of the content of Cr is determined from the viewpointof securing sufficient strength after quenching. Under a condition ofsecuring satisfactory cooling rate in quenching treatment, the contentof Cr is preferably from 0.05 to 0.3% by mass. When the strength afterquenching is strictly required even under varied cooling rate in thequenching treatment, the Cr content is preferably from 0.8 to 1.5% bymass.

The content of Mo is preferably from 0.05 to 0.5% by mass.

The present invention further provides a method for manufacturing highcarbon hot-rolled steel sheet, having the steps of hot-rolling, primarycooling, holding, coiling, acid washing, and annealing.

The hot-rolling step applies hot-rolling to a steel consistingessentially of, in terms of percentages of mass, 0.10 to 0.70% C, 2.0%or less Si, 0.20 to 2.0% Mn, 0.03% or less P, 0.03% or less S, 0.1% orless Sol.Al, 0.01% or less N, and balance of Fe and inevitableimpurities, at finishing temperatures of (Ar₃ transformation point −10°C.) or above.

The steel may further contain at least one element selected from thegroup consisting of, in terms of percentages of mass, 0.05 to 1.5% Crand 0.01 to 0.5% Mo.

The steel may further contain at least one element selected from thegroup consisting of, in terms of percentages of mass, 0.005% or less B,1.0% or less Cu, 1.0% or less Ni, and 0.5% or less W.

The steel may further contain at least one element selected from thegroup consisting of, in terms of percentages of mass, 0.05 to 1.5% Crand 0.01 to 0.5% Mo, and further at least one element selected from thegroup consisting of, in terms of percentages of mass, 0.005% or less B,1.0% or less Cu, 1.0% or less Ni, and 0.5% or less W.

The steel may further contain at least one element selected from thegroup consisting of, in terms of percentages of mass, 0.5% or less Ti,0.5% or less Nb, 0.5% or less V, and 0.5% or less Zr.

The primary cooling step is primary cooling of a hot-rolled steel sheetdown to the cooling termination temperatures ranging from 450° C. to600° C. at cooling rates of higher than 120° C./sec. The upper limit ofthe cooling rate is preferably 700° C./sec from the point of facilitycapacity.

The holding step is to hold the cooled hot-rolled steel sheet in atemperature range from 450° C. to 650° C. by the secondary cooling untilcoiling.

The coiling step is to coil the cooled hot-rolled steel sheet at coilingtemperatures of 600° C. or below. The coiling temperature is preferablyin a range from 200° C. to 600° C.

The acid washing step is to apply acid washing to the coiled hot-rolledsteel sheet.

The annealing step is to anneal the hot-rolled steel sheet after theacid washing at temperatures ranging from 680° C. to Ac₁ transformationpoint.

The percentage indicating the composition of steel, referred to herein,is percentage by mass.

The present invention suppresses the generation of voids at punched endface during punching, and delays the growth of cracks during burring. Asa result, the present invention provides a high carbon hot-rolled steelsheet having 440 MPa or higher tensile strength and extremely excellentductility and stretch-flange formability. By applying the high carbonhot-rolled steel sheet having excellent ductility and stretch-flangeformability according to the present invention to highly durable partssuch as transmission parts represented by gear, advanced level offorming is attained in the forming step, which provides high productquality and allows manufacturing the parts at low cost with decreasednumber of manufacturing steps. Also for the parts of driving system, theintegrally formed components are requested to have increased strength inthe non-heat treating parts for attaining higher durability and lighterweight, thus the steel sheets as the starting material are requested tohave 440 MPa class tensile strength (TS). The high carbon hot-rolledsteel sheet according to the present invention is useful in thisrespect.

DESCRIPTION OF THE EMBODIMENTS

The high carbon hot-rolled steel sheet according to the presentinvention consists essentially of, in terms of percentages of mass, 0.1to 0.7% C, 2.0% or less Si, 0.2 to 2.0% Mn, 0.03% or less P, 0.03% orless S, 0.1% or less Sol.Al, 0.01% or less N, and balance of Fe andinevitable impurities, and has a structure of ferrite having 6 μm orsmaller average grain size and carbide having 0.10 μm or more and lessthan 1.2 μm of average grain size, wherein the volume ratio of thecarbide having 2.0 μm or larger grain size is 10% or less, and thevolume ratio of the ferrite containing no carbide is 5% or less. Theabove specification of the steel sheet is most important parameter ofthe present invention. With thus specified chemical composition,metallic structure (average grain size of ferrite), shape of carbide(volume ratio of carbide having 2.0 μm or larger average grain size),and dispersion state of carbide (volume ratio of carbide-free ferritegrains), and by satisfying all of these specifications, a high carbonhot-rolled steel sheet having excellent ductility and stretch-flangeformability is obtained.

The high carbon hot-rolled steel sheet according to the presentinvention may further contain one or both of, in terms of percentage bymass, 0.05 to 1.5% C and 0.01 to 0.5% Mo, may further contain one ormore of, in terms of percentage by mass, 0.005% or less B, 1.0% or lessCu, 1.0% or less Ni, and 0.5% or less W, and may further contain one ormore of, in terms of percentage by mass, 0.5% or less Ti, 0.5% or lessNb, 0.5% or less V, and 0.5% or less Zr.

The high carbon hot-rolled steel sheet can be manufactured by the stepsof: hot-rolling the steel at finishing temperatures of (Ar₃transformation point −10° C.) or above; applying primary cooling to thehot-rolled steel sheet down to cooling termination temperatures rangingfrom 450° C. to 600° C. at cooling rates of higher than 120° C./sec;applying secondary cooling to hold the primarily cooled hot-rolled steelsheet in a temperature range from 450° C. to 650° C. until coiling;coiling the cooled hot-rolled steel sheet at coiling temperatures of600° C. or below; applying acid washing to the coiled hot-rolled steelsheet; and annealing the acid-washed hot-rolled steel sheet at annealingtemperatures ranging from 680° C. to Ac₁ transformation point. Theobject of the invention is attained by totally controlling theconditions of, after the hot-rolling, primary cooling, secondarycooling, coiling, and annealing.

The present invention is described in more detail in the following.

The reasons to limit the chemical composition of the steel according tothe present invention are described below.

C: 0.1 to 0.7%

Carbon is an important element that forms carbide and provides hardnessafter quenching. However, the C content of less than 0.1% causesconspicuous formation of proeutectoid ferrite in the structure after thehot-rolling, which results in uneven carbide distribution. In such acase, strength sufficient for structural machine parts cannot beobtained even after quenching. On the other hand, the C contentexceeding 0.7% results in insufficient working property, giving lowstretch-flange formability and ductility. In such a case, the steelsheet after the hot-rolling shows high hardness and becomes brittle sothat the strength after quenching saturates. Therefore, the C content isspecified to a range from 0.1 to 0.7%. From the point of securingsufficient strength after quenching, the C content is preferably 0.2,%or more, and from the point of handling of steel sheet on and aftercoiling, the C content is preferably 0.6% or less. The C contentcondition is an important parameter of the present invention.

Si: 2.0% or Less

Since Si is an element to improve the quenching property and increasethe material strength by solid solution strengthening, the Si content ispreferably 0.005% or more. However, The Si content exceeding 2.0%facilitates formation of proeutectoid ferrite and increases the ferritegrains substantially free from carbide, thereby deteriorating thestretch-flange formability. Furthermore, Si has a tendency ofgraphitizing carbide and likely hinders quenching property.Consequently, the Si content is specified to 2.0% or less, preferably0.02% or more from the point of securing strength after annealing, andpreferably 0.5% or less from the point of surface property.

Mn: 0.2 to 2.0%

Similar with Si, Mn is an element to improve the quenching property andto increase the material strength by solid solution strengthening.Manganese is also an important element which fixes S as MnS and preventshot cracking of slab. However, the Mn content of less than 0.2% reducesthese effects, and enhances the formation of proeutectoid ferrite togenerate coarse ferrite grains, and further significantly deterioratesthe quenching property. The Mn content exceeding 2.0% allows significantformation of manganese band which is a segregation zone, though a wantedtensile strength is attained, thereby deteriorating the stretch-flangeformability and the elongation. Accordingly, the Mn content is specifiedto a range from 0.20% to 2.0%, and preferably 1.0% or less from theviewpoint of stretch-flange formability and deterioration in elongationcaused by the formation of manganese band.

P: 0.03% or Less

Phosphorus is an element to be reduced because P is segregated in grainboundaries to decrease the toughness. Since, however, the P content isacceptable up to 0.03%, the P content is specified to 0.03% or less.

S: 0.03% or Less

Sulfur is an element to be reduced because S forms MnS with Mn todeteriorate the stretch-flange formability. Since, however, the Scontent is acceptable up to 0.03%, the S content is specified to 0.03%or less.

sol.Al: 0.1% or less

Aluminum is added in the steel making stage as an acid-eliminating agentto improve the cleanliness of steel. Normally Al is contained in thesteel in an amount of 0.005% or more as sol.Al. An Al content exceeding0.1% as sol.Al results in the saturation of the cleanliness improvingeffect, thereby increasing the cost. In addition, excess Al results inlarge amount of AlN precipitate to deteriorate the quenching property.Therefore, the sol.Al content is specified to 0.1% or less, preferably0.08% or less.

N: 0.01% or Less

Since excess N deteriorates the ductility, the N addition is specifiedto 0.01% or less.

The steel sheet according to the present invention achieves theobjective characteristics with the above essential adding elements.Depending on the wanted characteristics, however, one or both of Cr andMo may be added.

Cr: 0.05 to 1.5%

Chromium is an important element to suppress the formation ofproeutectoid ferrite during cooling step after the hot-rolling, thus toimprove the stretch-flange formability and improve the quenchingproperty. However, the Cr content less than 0.05% cannot attainsatisfactory effect. Furthermore, the Cr content exceeding 1.5%saturates the effect to suppress the formation of proeutectoid ferriteand increases the cost, though the quenching property is improved.Accordingly, when Cr is added, the Cr content is specified to a rangefrom 0.05 to 1.5%. Preferably, from the point of securing sufficientstrength after quenching, the Cr content is in a range from 0.05 to 0.3%under a condition that a satisfactory cooling rate is assured atquenching, and from 0.8 to 1.5% when a strict strength condition isrequested after quenching even under varied cooling rate at quenching.

Mo: 0.01 to 0.5%

Molybdenum is an important element to suppress the formation ofproeutectoid ferrite during the cooling step after the hot-rolling, thusto improve the stretch-flange formability and improve the quenchingproperty. However, the Mo content of less than 0.01% cannot attainsatisfactory effect. On the other hand, the Mo content exceeding 0.5%saturates the effect to suppress the formation of proeutectoid ferriteand increases the cost, though the quenching property is improved.Accordingly, when Mo is added, the Mo content is specified to a rangefrom 0.01 to 0.5%, and preferably 0.05% or more from the point ofsecuring sufficient strength after quenching.

The steel according to the present invention may further contain, addingto the above adding elements, one or more of B, Cu, Ni, and W, at need,to suppress the formation of proeutectoid ferrite during hot-rolling andcooling and to improve the quenching property. In such a case, less than0.0001% B, and less than 0.01% for each of Cu, Ni, and W cannot fullyattain the added effect. On the other hand, the added quantity exceeding0.005% B, 1.0% Cu, 1.0% Ni, and 0.5% W saturates the added affect, andincreases the cost. Consequently, on adding these elements, thespecified content is 0.0001 to 0.005% B, 0.01 to 1.0% Cu, 0.01 to 1.0%Ni, and 0.01 to 0.5% W. Boron, however, may form a compound with N inthe steel to fail in providing the effect of B itself. Therefore, theelement to be added for suppressing the formation of proeutectoidferrite during hot-rolling and cooling and for improving the quenchingproperty is preferably selected by one or more among the elements of Cu,Ni, and W. In that case, preferable adding amount of the respectiveelements is similar with that given above.

The steel according to the present invention may further contain, addingto the above adding elements, one or more of Ti, Nb, V, and Zr forassuring 440 MPa or higher tensile strength by refining the ferritegrains. In that case, each content less than 0.001% cannot obtainsufficient effect of addition. On the other hand, each content exceeding0.5% saturates the adding effect and increases the cost. Therefore, ifthese elements are added, the content of each one is specified to arange from 0.001 to 0.5%.

The balance to the above composition is Fe and inevitable impurities.

During the manufacturing process, various elements such as Sn and Pb mayenter as impurities. Those kinds of impurities, however, do notinfluence the effect of the present invention.

The following is the description of the present invention in terms ofmetallic structure (average grain size of ferrite), shape of carbide(average grain size of carbide and volume ratio of carbide having 2.0 μmor larger average grain size), and dispersion state of carbide (volumeratio of carbide-free ferrite grains). These conditions are importantparameters to obtain the high carbon hot-rolled steel sheet havingexcellent ductility and stretch-flange formability, and the effect ofthe present invention cannot be attained if any of these conditions isnot satisfied, or the effect of the present invention is attained onlyafter satisfying all of these conditions.

Average Ferrite Grain Size: 6 μm or Smaller

The average ferrite grain size is an important parameter governing thestretch-flange formability and the material strength. By refining theferrite grains, the strength is increased without deteriorating thestretch-flange formability. More specifically, average ferrite grainsizes of 6 μm or smaller provide excellent ductility and stretch-flangeformability while securing 440 MPa or higher tensile strength of thematerial. The average ferrite grain size can be controlled by theprimary cooling termination temperature, the secondary cooling holdingtemperature, and the coiling temperature, after hot-rolling, which aredescribed below.

Average Carbide Grain Size: 0.10 μm or Larger and Smaller Than 1.2 μm

The average carbide grain size significantly influences the workingproperties in general and the void formation during burring. Thus theaverage carbide grain size is an important parameter of the presentinvention. Although smaller carbide grain sizes suppress more the voidformation, average carbide grain size of smaller than 0.10 μmdeteriorates the ductility with the increase in hardness, therebydeteriorating the stretch-flange formability. On the other hand,increased average carbide grain size generally improves the workingproperty. The size exceeding 1.2 μm, however, leads to void formationduring burring to deteriorate the stretch-flange formability, andfurther the decrease in the local ductility causes the deterioration ofductility. Consequently, the average carbide grain size is specified toa range from 0.10 μm or larger and smaller than 1.2 μm. As describedbelow, the average carbide grain size can be controlled by themanufacturing conditions, specifically by the primary coolingtermination temperature, the coiling temperature, and the annealingtemperature.

Volume Ratio of Carbide Having 2.0 μm or Larger Grain Size: 10% or Less

During general working process and burring step, voids predominantlyoccur in the vicinity of coarse carbide. Accordingly, carbide has to beemphasized to control the average grain size and to reduce the volumeratio of coarse carbide grains, and they are also important parametersof the present invention. Even when the average carbide grain size is ina range from 0.10 μm or larger and smaller than 1.2 μm, the existence ofmore than 10% volume ratio of coarse carbide grains at or larger than2.0 μmin size deteriorates the stretch-flange formability caused by thegeneration of voids during burring, thereby decreasing the localductility to result in the deterioration of ductility. Consequently, thevolume ratio of the carbide having 2.0 μm or larger grain size isspecified to 10% or less. As described below, the carbide grain size canbe controlled by the primary cooling termination temperature, thesecondary cooling holding temperature, the coiling temperature, and theannealing temperature.

Volume Ratio of Carbide-Free Ferrite Grain Size: 5% or Less

Uniform dispersion of carbide relaxes the stress concentration on apunched end face during burring, thereby suppressing the void formation.In this regard, it is important to control the volume ratio ofcarbide-free ferrite grains. By controlling the volume ratio ofcarbide-free ferrite grains to 5% or less, the effect similar with thestate of uniform dispersion of carbide is attained, and thestretch-flange formability is significantly improved. In addition, localductility is improved, which then significantly improves the ductility.The term “carbide-free” referred to herein signifies that no carbide isdetected in an ordinary metal structure observation (with an opticalmicroscope). That type of ferrite grains forms a zone appeared as theproeutectoid ferrite after hot-rolling, where substantially no carbideis observed within grain even after the annealing. As described below,the state of carbide dispersion can be controlled by the manufacturingconditions, specifically by the finishing temperature, the cooling rateduring cooling after the rolling, the cooling termination temperature,and the coiling temperature.

The following is the description about the manufacturing method for highcarbon hot-rolled steel sheet having excellent ductility andstretch-flange formability according to the present invention.

The high carbon hot-rolled steel sheet according to the presentinvention is obtained by the steps of: hot-rolling a steel prepared tohave the above range of chemical composition at finishing temperaturesof (Ar₃ transformation point −10° C.) or above; applying primary coolingto the hot-rolled steel sheet down to cooling termination temperaturesranging from 450° C. to 600° C. at cooling rates of higher than 120°C./sec; applying secondary cooling to hold the primarily cooledhot-rolled steel sheet in a temperature range from 450° C. to 650° C.until coiling; coiling the cooled hot-rolled steel sheet at coilingtemperatures of 600° C. or below; applying acid washing to the coiledhot-rolled steel sheet; and annealing the acid-washed hot-rolled steelsheet at annealing temperatures ranging from 680° C. to Ac₁transformation point. The detail of the respective steps is describedbelow.

Finishing Temperature: Hot-Rolling at (Ar₃ Transformation Point −10° C.)or Above

Finishing temperature of hot-rolling below (Ar₃ transformation point−10° C.) enhances the ferrite transformation in a part, which increasesthe ferrite grains to deteriorate the ductility and the stretch-flangeformability. Therefore, the finish-rolling is done at finishingtemperatures of (Ar₃ transformation point −10° C.) or above. Thecondition assures uniform structure and improves the ductility and thestretch-flange formability.

Cooling Rate: Primary Cooling at Rates of Higher Than 120° C./sec

According to the present invention, rapid cooling (primary cooling) isadopted at cooling rates of higher than 120° C./sec after hot-rolling toreduce the volume ratio of proeutectoid ferrite after transformation.Gradual cooling results in a low super cooling degree of austenite,leading to the formation of proeutectoid ferrite. In particular, 0.120°C./sec or smaller cooling rate gives conspicuous formation ofproeutectoid ferrite, thereby resulting in the carbide-free ferritegrains exceeding 5% to deteriorate the ductility and the stretch-flangeformability. Accordingly, the cooling rate after hot-rolling isspecified to higher than 120° C./sec.

It is preferable to begin the primary cooling after the finish-rollingwithin a period of from more than 0.1 sec and less than 1.0 sec. Thecondition provides finer ferrite grains and precipitates such aspearlite after the transformation, thus further improving the workingproperty.

Cooling Termination Temperature: 450° C. to 600° C.

High cooling termination temperature in the primary cooling causesproeutectoid ferrite formation and increase in the lamella spacing ofpearlite. As a result, fine carbide cannot be obtained after theannealing, and the ductility and the stretch-flange formability aredeteriorated. Particularly when the cooling termination temperature ishigher than 600° C., the carbide-free ferrite grains increase to morethan 5%, which deteriorates the ductility and the stretch-flangeformability. Therefore, the cooling termination temperature afterrolling is specified to 600° C. or below. Lower than 450° C. of coolingtermination temperature cannot obtain the equiaxed ferrite grains, anddeteriorates the working property. Therefore, the cooling terminationtemperature is specified to 450° C. or above.

Secondary Cooling From the Primary Cooling Termination to the Coiling:Holding at Temperatures in a Range From 450° C. to 650° C.

For the case of high carbon steel sheets, the steel sheet temperatureincreases after the primary cooling termination, in some cases,accompanied by the proeutectoid ferrite transformation, the pearlitetransformation, and the bainite transformation. Thus, even if theprimary cooling termination temperature is lower than 600° C., when thetemperature in the course from the primary cooling termination to thecoiling is higher than 650° C., the proeutectoid ferrite is formed, thelamella spacing of pearlite increases, and the carbide in pearlitebecomes coarse. As a result, the fine carbide cannot be obtained afterthe annealing, and the volume ratio of carbide having 2.0 μm or largergrain size exceeds 10%, thereby deteriorating the ductility and thestretch-flange formability. If the temperature in the course from theprimary cooling termination to the coiling is lower than 450° C., theequiaxed ferrite cannot be obtained to deteriorate the working property,in some cases. Therefore, it is important to control the temperature inthe course from the secondary cooling to the coiling. By holding thematerial between the secondary cooling step and the coiling step totemperatures ranging from 450° C. to 650° C., the deterioration ofductility, of stretch-flange formability, and of working property can beprevented. The secondary cooling may be done by laminar cooling or thelike.

Regarding the holding time from the primary cooling termination to thecoiling, short in the time induces the generation of transformation heatafter coiling, which makes the steel sheet temperature controlimpossible and generates coil crushing. Therefore, the holding time ispreferably 5 seconds or more for completing the transformation untilcoiling, and preferably 60 seconds or less because excess holding timesignificantly deteriorates the operability.

Coiling Temperature: 600° C. or below

Higher coiling temperature increases more the lamella spacing ofpearlite. Thus, the carbide becomes coarse after the annealing. When thecoiling temperature exceeds 600° C., the ductility and thestretch-flange formability deteriorate. Consequently, the coilingtemperature is specified to 600° C. or below. Although the lower limitof the coiling temperature is not specifically defined, 200° C. or aboveis preferred because lower temperature induces more the deterioration ofsteel sheet shape.

Annealing Temperature: 680° C. to Ac₁ Transformation Point

After applying acid washing to the hot-rolled steel sheet, annealing isgiven for spheroidizing the carbide. The annealing temperature lowerthan 680° C. results in insufficient spheroidization of carbide or informing carbide having smaller than 0.1 μm of average grain size, whichdeteriorates the stretch-flange formability. In addition, no equiaxedferrite is obtained, and the working property and the ductility aredeteriorated. On the other hand, annealing temperature exceeding the Ac₁transformation point causes austenite formation in a part, which againgenerates pearlite during cooling, thereby also deteriorating thestretch-flange formability and the ductility. Consequently, theannealing temperature is specified to a range from 680° C. to Ac₁transformation point.

For the composition preparation of the high carbon steel according tothe present invention, either a converter or an electric furnace can beapplied. The high carbon steel after the composition preparation isformed in a steel slab by block formation—block rolling or by continuouscasting. The steel slab is subjected to hot-rolling. The slab heatingtemperature is preferably 1280° C. or below to avoid deterioration ofthe surface state caused by scaling. The continuously cast slab may besent, in as-cast state, to direct-feed rolling in which the slab isrolled under heating to prevent temperature reduction. Furthermore,finish-rolling may be given during the hot-rolling step eliminating therough-rolling. Alternatively, to secure the finishing temperature, therolled material may be heated with a heating means such as bar heaterduring the hot-rolling. Also in order to accelerate spheroidization orto reduce the hardness, the coiled steel sheet may be held to thetemperature with a gradual cooling cover or other means.

The annealing after hot-rolling may be conducted by box annealing orcontinuous annealing. Temper rolling is succeedingly executed at need.Since the temper rolling does not influence the quenching property, thecondition of temper rolling is not specifically limited.

The above procedure provides a high carbon hot-rolled steel sheet havingexcellent ductility and stretch-flange formability. A presumable reasonthat the high carbon hot-rolled steel sheet according to the presentinvention has the excellent ductility and stretch-flange formability isthe following. The stretch-flange formability is significantly affectedby the internal structure of punched end face zone. It was confirmedthat, particularly for the case of large amount of carbide-free ferritegrains (the proeutectoid ferrite after the hot-rolling), cracks aregenerated from the grain boundary with the spheroidal structure zone.When the behavior of microstructure is observed, the void formationcaused by the stress concentration becomes stronger at the interface ofcarbide after the punching. The stress concentration is enhanced in astate of increased size of carbide grains and increased quantity ofcarbide-free ferrite grains. On burring, these voids are connected eachother to form cracks. Further by controlling the ferrite grain size, theelongation stably increases. From the above phenomena, it is possible toreduce stress concentration, to reduce void generation, thus to provideexcellent ductility and stretch-flange formability through the controlof chemical composition, metallic structure (average ferrite grainsize), carbide shape (volume ratio of carbide having 2.0 μm or largeraverage grain size), and dispersed state of carbide (volume ratio ofcarbide-free ferrite grains).

EXAMPLE 1

Continuously cast slabs of steels having the respective chemicalcompositions given in Table 1 as the steel Nos. A to R were heated to1250° C., then were subjected to hot-rolling and annealing under therespective conditions given in Table 2 to prepare steel sheets having5.0 mm in thickness. The steel sheet Nos. 1 to 18 are the example steelsprepared under the manufacturing conditions within the range of thepresent invention, and the steel Nos. 19 to 32 are the comparativeexample steels prepared under the manufacturing conditions outside therange of the present invention.

Samples were cut from thus prepared respective steel sheets, and weresubjected to measurements of ferrite grain size, average carbide grainsize, volume ratio of carbide having 2.0 2 m or larger grain size,volume ratio of carbide-free ferrite grains, hardness, andstretch-flange formability (burring ratio), and further to tensile test.The results are given in Table 3. Method and condition of each test andmeasurement are the following.

(1) Determination of Ferrite Grain Size, Average Carbide Grain Size,Volume Ratio of Carbide Having 2.0 μm or Larger Grain Size, and VolumeRatio of Carbide-Free Ferrite Grains

A cross section along the thickness of a sample sheet was polished,etched, and photographed by a scanning electron microscope to observethe microstructure within an area of 0.01 mm²; The determination wasgiven on the ferrite grain size, the average carbide grain size, thevolume ratio of carbide having 2.0 μm or larger grain size, and thevolume ratio of carbide-free ferrite grains.

(2) Determination of Hardness

The surface hardness of steel sheet was determined in accordance withJIS Z2245. Average of n=5 data was derived.

(3) Determination of Stretch-Flange Formability

A sample was punched with a punching tool having a punch diameter ofd₀=10 mm and a die diameter of 12 mm (clearance 20%), and was subjectedto a hole-expanding test. The hole-expanding test was executed by thepush-up method with a cylindrical flat-bottomed punch (50 mmf, 8R)),then a hole diameter db was measured when a crack was generated acrossthe thickness of the sheet. The hole-expanding ratio λ(%) defined by thefollowing formula was derived.λ=100×(db−d ₀)/d ₀  (1)(4) Tensile Test

A JIS No. 5 sheet was cut along the direction of 90° (C direction) tothe rolling direction, and was subjected to tensile test with a testingspeed of 10 mm/min to determine the tensile strength and the elongation.

The present invention places the target values of: 440 MPa or highertensile strength TS; 35% or higher elongation for a steel containing0.10% or more and less than 0.40% C; 30% or higher elongation for asteel containing 0.40 to 0.70% C; 70% or higher hole-expanding ratio λfor a steel containing 0.10% or more and less than 0.40% C (5.0 mm ofsheet thickness); and 40% or higher hole expanding ratio λ for a steelcontaining 0.40 to 0.70% C (5.0 mm of sheet thickness).

Table 3 shows that the example steel sheet Nos. 1 to 18 of the presentinvention gave 440 MPa or higher tensile strength (TS), with highhole-expanding ratio λ, thus providing excellent stretch-flangeformability and elongation.

In contrast, the steel sheet Nos. 19 to 32 are the comparative examplesteels which were prepared under the manufacturing conditions outsidethe range of the present invention. The steel sheet Nos. 19, 20, 22, 23,and 24 gave the ferrite grain size larger than 6 μm so that theirtensile strengths were below 440 MPa. The steel sheet Nos. 30 and 31gave the average carbide grain size larger than 1.2 μm so that theirvolume ratio of carbide having larger than 2 μm of the grain sizeexceeded 10%, and further their volume ratio of carbide-free ferriteexceeded 5%, thus the hole-expanding ratio λ was low, and thestretch-flange formability was poor. The steel sheet Nos. 21, 25, 28,and 32 gave smaller than 0.1 μm of average carbide grain size toincrease the strength so that the hole expanding ratio λ and theelongation were low compared with the target values, and the elongationand the stretch-flange formability were poor. The steel sheet Nos. 27and 29 gave larger than 5% in the volume ratio of carbide-free ferriteso that the hole expanding ratio λ and the elongation were low comparedwith the target values, and the elongation and the stretch-flangeformability were poor. The steel sheet No. 26 gave more than 10% of thevolume ratio of carbide having 2.0 μm or larger grain size, though theaverage carbide grain size was in a range from 0.10 μm or larger andsmaller than 1.2 μm, thus the hole expanding ratio λ and the elongationwere low compared with the target values, and the stretch-flangeformability and the elongation were poor. TABLE 1 Steel No. C Si Mn P Ssol. Al N Other A 0.15 0.22 0.72 0.009 0.005 0.020 0.0038 Cr: 1.0, Mo:0.16 B 0.23 0.20 0.80 0.010 0.009 0.031 0.0030 — C 0.35 0.21 0.76 0.0140.005 0.028 0.0034 — D 0.35 0.20 0.75 0.012 0.004 0.035 0.0036 Cr: 1.0,Mo: 0.16 E 0.49 0.18 0.75 0.011 0.008 0.030 0.0035 — F 0.64 0.22 0.730.012 0.010 0.021 0.0036 — G 0.26 0.03 0.45 0.015 0.003 0.040 0.0050 Cr:0.28 H 0.26 0.03 0.45 0.015 0.003 0.040 0.0050 Mo: 0.30 I 0.47 0.18 0.750.011 0.008 0.030 0.0035 Cr: 0.15 J 0.58 0.20 0.74 0.015 0.010 0.0210.0038 Cr: 0.06 K 0.35 0.21 0.76 0.013 0.005 0.028 0.0034 Cr: 0.18 L0.35 0.45 0.76 0.013 0.005 0.028 0.0034 Mo: 0.06 M 0.37 0.03 0.75 0.0140.004 0.028 0.0034 Cr: 0.28, Mo: 0.30 N 0.35 0.18 0.25 0.014 0.005 0.0280.0034 Mo: 0.15 O 0.35 0.18 0.95 0.014 0.005 0.028 0.0034 Cr: 0.06, Mo:0.06 P 0.35 0.20 0.75 0.014 0.004 0.031 0.0032 Cr: 0.06, B: 0.0022, Cu:0.2, Ni: 0.6, W: 0.05 Q 0.34 0.21 0.75 0.013 0.004 0.032 0.0034 Cr:0.25, Ti: 0.005, Nb: 0.008, V: 0.01, Zr: 0.01 R 0.34 0.21 0.73 0.0130.004 0.030 0.0038 Cr: 0.06, Mo: 0.06, Cu: 0.08, Ni: 0.02, Ti: 0.02, V:0.05

TABLE 2 Primary Primary Rolling cooling Primary cooling Range of holdingSteel termination starting cooling termination temperature Coiling sheetSteel temperature time rate temperature in the secondary coolingtemperature Annealing No. No. (° C.) (sec) (° C./sec) (° C.) until thecoiling (° C.) (° C.) condition Remark 1 A Ar3 + 30° C. 0.5 220 590550˜590 550 680° C. × 40 hr Example 2 B Ar3 + 30° C. 1.2 230 590 570˜620580 680° C. × 40 hr Example 3 C Ar3 + 20° C. 1.0 210 560 480˜550 540680° C. × 40 hr Example 4 D Ar3 + 20° C. 1.0 200 550 490˜530 540 680° C.× 40 hr Example 5 E Ar3 + 30° C. 1.2 200 570 520˜630 550 710° C. × 40 hrExample 6 F Ar3 + 40° C. 0.4 200 580 580˜640 560 700° C. × 40 hr Example7 G Ar3 + 20° C. 1.1 210 590 580˜630 560 680° C. × 40 hr Example 8 HAr3 + 20° C. 1.1 220 580 580˜620 570 680° C. × 40 hr Example 9 I Ar3 +30° C. 1.2 210 560 530˜630 560 680° C. × 40 hr Example 10 J Ar3 + 20° C.1.1 200 570 540˜620 550 680° C. × 40 hr Example 11 K Ar3 + 20° C. 1.0210 560 480˜550 550 680° C. × 40 hr Example 12 L Ar3 + 20° C. 1.0 210570 480˜570 570 680° C. × 40 hr Example 13 M Ar3 + 20° C. 1.0 210 560480˜550 560 680° C. × 40 hr Example 14 N Ar3 + 20° C. 1.0 210 560480˜540 550 680° C. × 40 hr Example 15 O Ar3 + 20° C. 1.0 210 570480˜550 560 680° C. × 40 hr Example 16 P Ar3 + 20° C. 1.0 210 560490˜580 560 680° C. × 40 hr Example 17 Q Ar3 + 20° C. 1.0 210 560500˜570 560 680° C. × 40 hr Example 18 R Ar3 + 20° C. 1.0 210 560500˜570 560 680° C. × 40 hr Example 19 A Ar3 + 30° C. 0.5 180 680620˜650 600 680° C. × 40 hr Comparative Example 20 A Ar3 − 40° C. 1.2180 590 580˜630 590 680° C. × 40 hr Comparative Example 21 A Ar3 + 10°C. 0.5 280 430 420˜500 500 660° C. × 40 hr Comparative Example 22 BAr3 + 30° C. 1.2 210 630 580˜660 580 680° C. × 40 hr Comparative Example23 B Ar3 − 40° C. 0.7 160 630 560˜620 570 700° C. × 40 hr ComparativeExample 24 B Ar3 + 20° C. 1.2  80 610 550˜600 540 680° C. × 40 hrComparative Example 25 C Ar3 + 30° C. 0.8 220 580 470˜550 460 600° C. ×40 hr Comparative Example 26 C Ar3 + 20° C. 1.0 210 580 550˜680 600 680°C. × 40 hr Comparative Example 27 D Ar3 − 30° C. 1.2 160 590 580˜640 590680° C. × 40 hr Comparative Example 28 D Ar3 + 20° C. 0.5 280 420410˜510 500 660° C. × 40 hr Comparative Example 29 E Ar3 − 30° C. 1.2160 580 550˜630 520 700° C. × 40 hr Comparative Example 30 E Ar3 + 30°C. 0.7 200 660 610˜650 600 700° C. × 40 hr Comparative Example 31 FAr3 + 20° C. 1.0 180 640 600˜650 640 700° C. ×40 hr Comparative Example32 F Ar3 + 10° C. 0.6 220 610 540˜610 560 640° C. × 40 hr ComparativeExample

TABLE 3 Average Average Volume ratio of Volume ratio Steel ferritecarbide carbide larger than of Tensile sheet Steel grain grain 2 μm ingrain carbide-free Hardness Hole-expanding strength Elongation No. No.size (μm) size (μm) size (%) ferrite (HRB) ratio λ (%) (MPa) (%) Remark1 A 5.8 0.75 6 5 73 148 440 43 Example 2 B 5.5 0.88 8 5 73 150 445 42Example 3 C 3.6 0.59 4 3 79 80 490 38 Example 4 D 3.2 0.40 2 3 80 75 50036 Example 5 E 2.9 0.47 3 2 86 56 560 32 Example 6 F 1.9 0.36 2 1 88 45590 31 Example 7 G 5.0 0.65 7 4 75 90 470 40 Example 8 H 4.8 0.63 6 4 7689 480 40 Example 9 I 3.0 0.50 3 2 85 60 550 33 Example 10 J 2.5 0.41 21 87 50 580 31 Example 11 K 3.6 0.57 3 3 79 79 490 38 Example 12 L 3.60.58 4 4 80 78 500 37 Example 13 M 3.6 0.59 4 3 78 81 480 39 Example 14N 3.6 0.59 4 3 79 80 490 38 Example 15 O 3.6 0.59 4 3 79 79 490 38Example 16 P 3.5 0.58 4 3 79 79 490 38 Example 17 Q 3.2 0.58 4 3 80 78500 37 Example 18 R 3.2 0.59 4 3 79 80 490 38 Example 19 A 10.8 1.44 2530 70 98 410 42 Comparative Example 20 A 6.8 0.90 9 20 72 118 435 40Comparative Example 21 A 3.5 0.05 0 1 84 38 535 33 Comparative Example22 B 6.5 0.94 11 8 72 138 430 40 Comparative Example 23 B 7.2 1.30 15 2668 75 400 41 Comparative Example 24 B 6.5 0.88 8 16 72 70 430 40Comparative Example 25 C 3.4 0.07 0 2 90 21 580 29 Comparative Example26 C 3.6 1.10 11 5 79 45 490 32 Comparative Example 27 D 5.2 0.64 5 1578 51 480 33 Comparative Example 28 D 2.1 0.04 0 0 92 20 600 27Comparative Example 29 E 3.0 0.68 6 18 82 19 520 28 Comparative Example30 E 5.2 1.39 22 15 80 20 500 29 Comparative Example 31 F 3.9 1.38 21 684 10 530 27 Comparative Example 32 F 3.0 0.08 1 6 89 11 580 25Comparative Example

1. A high carbon hot-rolled steel sheet consisting essentially of: interms of percentage of mass, 0.10 to 0.7% C, 2.0% or less Si, 0.20 to2.0% Mn, 0.03% or less P, 0.03% or less S, 0.1% or less Sol.Al, 0.01% orless N, and the balance being Fe and inevitable impurities; ferritehaving an average grain size of 6 μm or less and carbide having anaverage grain size of 0.10 μm or more and less than 1.2 μm; the carbidehaving a volume ratio of 10% or less regarding a grain size of 2.0 μm ormore; and the ferrite containing no carbide having a volume ratio of 5%or less.
 2. The high carbon hot-rolled steel sheet according to claim 1,further containing at least one element selected from the groupconsisting of, in terms of percentages of mass, 0.05 to 1.5% Cr and 0.01to 0.5% Mo.
 3. The high carbon hot-rolled steel sheet according to claim1, further containing at least one element selected from the groupconsisting of, in terms of percentages of mass, 0.005% or less B, 1.0%or less Cu, 1.0% or less Ni, and 0.5% or less W.
 4. The high carbonhot-rolled steel sheet according to claim 2, further containing at leastone element selected from the group consisting of, in terms ofpercentages of mass, 0.005% or less B, 1.0% or less Cu, 1.0% or less Ni,and 0.5% or less W.
 5. The high carbon hot-rolled steel sheet accordingclaim 1, further containing at least one element selected from the groupconsisting of, in terms of percentages of mass, 0.5% or less Ti, 0.5% orless Nb, 0.5% or less V, and 0.5% or less Zr.
 6. The high carbonhot-rolled steel sheet according to claim 1, wherein the content of Siis 0.005 to 2.0% by mass.
 7. The high carbon hot-rolled steel sheetaccording to claim 6, wherein the content of Si is 0.02 to 0.5% by mass.8. The high carbon hot-rolled steel sheet according to claim 1, whereinthe content of Mn is 0.2 to 1.0% by mass.
 9. The high carbon hot-rolledsteel sheet according to claim 2, wherein the content of Cr is 0.05 to0.3% by mass.
 10. The high carbon hot-rolled steel sheet according toclaim 2, wherein the content of Cr is 0.8 to 1.5% by mass.
 11. The highcarbon hot-rolled steel sheet according to claim 2, wherein the contentof Mo is from 0.05 to 0.5% by mass.
 12. A method for manufacturing ahigh carbon hot-rolled steel sheet, comprising the steps of: hot-rollinga steel consisting essentially of, in terms of percentages of mass, 0.10to 0.70% C, 2.0% or less Si, 0.20 to 2.0% Mn, 0.03% or less P, 0.03% orless S, 0.1% or less Sol.Al, 0.01% or less N, and the balance being Feand inevitable impurities, at finishing temperatures of (Ar₃transformation point −10° C.) or more; applying primary cooling to thehot-rolled steel sheet down to cooling termination temperatures rangingfrom 450° C. to 600° C. at cooling rates of more than 120° C./sec;applying secondary cooling to hold the primarily cooled hot-rolled steelsheet in a temperature range from 450° C. to 650° C. until coiling;coiling the cooled hot-rolled steel sheet at coiling temperatures of600° C. or less; applying acid washing to the coiled hot-rolled steelsheet; and annealing the acid-washed hot-rolled steel sheet at annealingtemperatures ranging from 680° C. to Ac₁ transformation point.
 13. Themethod according claim 12, wherein the cooling rate in the primarycooling step is in a range from 120 to 700° C./sec.
 14. The methodaccording to claim 12, wherein the coiling temperature is in a rangefrom 200° C. to 600° C.
 15. The high carbon hot-rolled steel sheetaccording claim 2, further containing at least one element selected fromthe group consisting of, in terms of percentages of mass, 0.5% or lessTi, 0.5% or less Nb, 0.5% or less V, and 0.5% or less Zr.
 16. The highcarbon hot-rolled steel sheet according claim 3, further containing atleast one element selected from the group consisting of, in terms ofpercentages of mass, 0.5% or less Ti, 0.5% or less Nb, 0.5% or less V,and 0.5% or less Zr.
 17. The high carbon hot-rolled steel sheetaccording claim 4, further containing at least one element selected fromthe group consisting of, in terms of percentages of mass, 0.5% or lessTi, 0.5% or less Nb, 0.5% or less V, and 0.5% or less Zr.